Synergistic Effect of Metal Oxide and Carbon Nanoparticles on the Thermal and Mechanical Properties of Polyimide Composite Films

In this paper, we report on novel polyimide (PI) nanocomposites filled with binary mixtures of metal oxide (either TiO2 or ZrO2) nanoparticles and nanocarbon (either carbon nanofibers (CNFs) or functionalized carbon nanotubes (CNTfs)). The structure and morphology of the materials obtained were comprehensively studied. An exhaustive investigation of their thermal and mechanical properties was performed. We revealed a synergistic effect of the nanoconstituents with regard to a number of functional characteristics of the PIs compared with single-filler nanocomposites, including thermal stability, stiffness (below and above glass transition temperature), yield point, and temperature of flowing. Moreover, the possibility of manipulating the properties of the materials by choosing a proper combination of the nanofillers was demonstrated. The results obtained can become a platform in the design of PI-based engineering materials with tailored characteristics capable of operating in extreme conditions.


Introduction
Over the years, polyimides (PIs) being excellent engineering materials, have widely been used in many fields of industry, e.g., aerospace, mechanical and chemical industries, microelectronics, as well as production of fuel cells and household appliances. Such a variety of applications of PIs is ensured by a great set of their outstanding characteristics, including superb thermal resistance and thermal stability, high mechanical strength and good insulating properties, superior chemical and radiation resistance, etc. [1][2][3][4][5]. However, the further development of science and technology demands more materials that can operate well under harsh conditions. To fulfil the increasing need for the high-performance materials, PI-based nanocomposites have been vigorously investigated [6][7][8][9][10][11].
Metal oxide (MO) nanoparticles is a huge class of nanofillers having received a great attention because of the superiority of the MO-containing polymer nanocomposites over unfilled materials [12][13][14][15]. The widespread use of metal oxides comes also from the possibility of tuning the physico-chemical and working properties of the nanocomposites by changing the synthetic techniques of the nanocomponents, since the synthesis conditions determine their size, morphology, surface stoichiometry (the number of oxygen vacancies), and functionality [15][16][17]. The good prospects of PI-based materials filled with a number of MO nanospecies, such as ZrO 2 , TiO 2 , NiO, CeO 2 , and ZnO are reported in the literature [18][19][20][21][22][23][24][25][26][27][28].
properties of PIs crucial for practical application is quite obvious. PIs being one of the foremost engineering polymers, their thermal and mechanical characteristics are of great interest. The modification of PI matrices with mixed fillers synergistically affecting a set of characteristics may well become a perspective tool in developing novel materials relevant to industrial usage.
In the present work, we address the problem of the development of advanced PI-based nanocomposites filled with binary mixtures of nanofillers, the latter being the combinations of either TiO 2 or ZrO 2 and either CNT or CNF. The following prerequisites determined the choice of the components: Both the oxides are non-toxic, low-cost, and widely used because they are able to augment the corrosion resistance of materials, possess a high dielectric constant, and can be used as pigment pieces for thermoregulating coatings in aerospace [55]. They are also characterized by high thermal and chemical stability. Moreover, these nanoparticles have demonstrated good application prospects in improving the functional properties of PIs [19,21,22,24,56,57]. We already discussed how the type, size, surface functionality, and concentration of these nanofillers can affect the thermal and mechanical properties of PI-based nanocomposites [58]. On the other hand, both CNT and CNF have long been used as reinforcing agents in PIs imparting greater rigidity and mechanical strength to the host polymers [59][60][61]. In this study we reveal the synergism of the aforementioned types of nanofillers in the context of simultaneous enhancement of key thermal and mechanical characteristics of PIs. We also investigate the structure and morphology of the nanocomposites developed.
Structural formulas of the elementary units of PIs used in this work are shown in Figure 1.

Synthesis and Treatment of Nanoparticles
The synthesis of ZrO2 nanoparticles with a high content of OH-groups on the surface was carried out in two stages. First, 25% ammonium hydroxide was added dropwise to a 1 M solution of ZrOCl2·8H2O until a white cheesy precipitate formed at pH = 14. Then, the ZrO(OH)2 precipitate was repeatedly washed with distilled water until a negative reaction to chloride ions occurred and it dried to a constant weight at 100 °C in a furnace. Next, 0.5 g of zirconium oxyhydroxide powder, ground in a mortar to a powdery state, was placed in a Teflon cell and was poured with 14.5 mL of distilled water. Then, the steel autoclave was hermetically sealed, heated at a rate of 5 °C/min, and kept at 250 °C for 4 h. The ZrO2 nanoparticles formed were removed from the autoclave, dried at 100 °C and cooled to room temperature naturally. The powder was characterized using a set of methods of physicochemical analysis [62]. The presence of the hydroxyl groups on the surface of the nanoparticles was confirmed by FTIR spectroscopy ( Figure S1). These nanoparticles were designated ZrO2(OH).
ZrO2 nanoparticles with a more hydrophobic surface were formed under solvothermal conditions. A suspension of zirconium acetylacetonate in toluene was subjected to the isothermal exposure for 72 h at a temperature of 250 °C and a pressure of 70 MPa. The formed particles were then repeatedly washed with ethanol, dried to a constant weight, and annealed in air at 500 °C for 2 h, or at 800 °C for 10 min, depending on the required crystallite size [63]. The nanoparticles of crystallite sizes 8 and 18 nm (according to X-ray phase analysis) were obtained and denoted by ZrO2(8nm) and ZrO2(18nm), correspondingly.
Solvothermal treatment under the same conditions was performed on a mixture of titanium butoxide with toluene in autoclaves. TiO2 nanoparticles with the anatase structure were formed [64], and annealed at 500 °C for 2 h.
In order to prevent the aggregation of CNT and improve their compatibility with the PI matrices [61,65] we modified their surfaces, exploiting the following technique: CNT were treated with boiling concentrated nitric acid for 36 h in a ratio CNT(g)/nitric acid(mL) 1:40. The powder obtained was washed with distilled water from nitric acid residues with the use of a centrifuge until pH = 7 was reached. The nanoparticles functionalized were then freeze-dried and denoted by 40% CNTf. The FTIR spectra of the pristine and processed CNT were registered to confirm the functionalization ( Figure S2).

Synthesis of R-BAPS Prepolymer
A 15% solution of PAA based on diamine BAPS and dianhydride R was obtained by the polycondensation method. First, the diamine was dissolved in NMP. Then, an

Synthesis and Treatment of Nanoparticles
The synthesis of ZrO 2 nanoparticles with a high content of OH-groups on the surface was carried out in two stages. First, 25% ammonium hydroxide was added dropwise to a 1 M solution of ZrOCl 2 ·8H 2 O until a white cheesy precipitate formed at pH = 14. Then, the ZrO(OH) 2 precipitate was repeatedly washed with distilled water until a negative reaction to chloride ions occurred and it dried to a constant weight at 100 • C in a furnace. Next, 0.5 g of zirconium oxyhydroxide powder, ground in a mortar to a powdery state, was placed in a Teflon cell and was poured with 14.5 mL of distilled water. Then, the steel autoclave was hermetically sealed, heated at a rate of 5 • C/min, and kept at 250 • C for 4 h. The ZrO 2 nanoparticles formed were removed from the autoclave, dried at 100 • C and cooled to room temperature naturally. The powder was characterized using a set of methods of physicochemical analysis [62]. The presence of the hydroxyl groups on the surface of the nanoparticles was confirmed by FTIR spectroscopy ( Figure S1). These nanoparticles were designated ZrO 2(OH) .
ZrO 2 nanoparticles with a more hydrophobic surface were formed under solvothermal conditions. A suspension of zirconium acetylacetonate in toluene was subjected to the isothermal exposure for 72 h at a temperature of 250 • C and a pressure of 70 MPa. The formed particles were then repeatedly washed with ethanol, dried to a constant weight, and annealed in air at 500 • C for 2 h, or at 800 • C for 10 min, depending on the required crystallite size [63]. The nanoparticles of crystallite sizes 8 and 18 nm (according to X-ray phase analysis) were obtained and denoted by ZrO 2 (8nm) and ZrO 2 (18nm), correspondingly.
Solvothermal treatment under the same conditions was performed on a mixture of titanium butoxide with toluene in autoclaves. TiO 2 nanoparticles with the anatase structure were formed [64], and annealed at 500 • C for 2 h.
In order to prevent the aggregation of CNT and improve their compatibility with the PI matrices [61,65] we modified their surfaces, exploiting the following technique: CNT were treated with boiling concentrated nitric acid for 36 h in a ratio CNT(g)/nitric acid(mL) 1:40. The powder obtained was washed with distilled water from nitric acid residues with the use of a centrifuge until pH = 7 was reached. The nanoparticles functionalized were then freeze-dried and denoted by 40% CNT f . The FTIR spectra of the pristine and processed CNT were registered to confirm the functionalization ( Figure S2).

Synthesis of R-BAPS Prepolymer
A 15% solution of PAA based on diamine BAPS and dianhydride R was obtained by the polycondensation method. First, the diamine was dissolved in NMP. Then, an equimolar amount of dianhydride was gradually added to the solution under continuous stirring and cooling in an ice bath. The resulting solution was stirred for 4 h at room temperature in argon flow. The formed prepolymer was filtered and degassed. Prior to the synthesis, all monomers were dried for 24 h at temperatures 20 • C below their melting points. NMP was distilled under vacuum just before using in the synthesis of R-BAPS polyimide.

Preparation of Pristine and Nanocomposite Films
All the PI-based compositions (both with single and binary nanofillers) were prepared employing a standard solution technique [66,67]. The proper amounts of the nanoparticles, either MO or nanocarbon or their mixture, were sonicated in NMP for 1 h and then blended with the corresponding PAA solutions. The sizes and concentrations of the MO nanoparticles were chosen based on our previous results to provide the diversity of thermal and mechanical behavior of PIs filled with these nanospecies [58]. The choice of the amounts of the nanocarbon species was also based on our previous tests (results are not presented) ensuring optimal properties of the single-filler nanocomposites. The nanocomposite solutions obtained were stirred for 24 h to form a quasi-homogeneous system. These compositions as well as the host PAA solutions were cast on glass supports and dried for 4 h at 80 • C. A gradual heating up to either 365 • C (for PMDA-ODA-based samples) or 300 • C (for R-BAPS-based samples) was performed at a rate of 3 • C/min in order to prevent bubbling of the solvent. The final curing at these temperatures was carried out for 30 min. A list of the samples obtained is provided in Table 1.

Characterization Techniques
We investigated the structural and morphological features of the pristine polymer as well as nanocomposite samples employing a number of techniques.
Optical images of the surfaces of the nanocomposite samples were obtained, employing an ADF PRO20 digital microscopy camera (ADF Optics CO, Ltd., Hangzhou, China).
Scanning electron microscopy (SEM) was carried out on the cryo-cleavages of the films using a SUPRA-55VP microscope (Carl Zeiss, Oberkochen, Germany) equipped with a secondary electron detector. The specimens were glued on the microscope holders and covered with a thin layer of platinum.
Atomic force microscopy (AFM) data were obtained using an SPM-9700HT scanning probe microscope (Shimadzu, Kyoto, Japan). The AFM images were captured in air at room temperature. The setup operated in a tapping mode using NSG-10 Silicon tips with curvature radius 5 nm. The images of 1024 × 1024 points were obtained.
FTIR spectra of unfilled and nanocomposite films were recorded on a Vertex 70 IR Fourier spectrometer (Bruker, Billerica, MA, USA) with the ATR (Attenuated Total Reflection) reflector (Pike Technologies, WI, USA) at room temperature in the range of 4000-400 cm -1 (number of scans 30) with a ZnSe working element. When registering the ATR spectra, a correction was made that includes the penetration depth depending on the wavelength. The equipment was also employed to confirm the presence of functional groups on the surfaces of ZrO 2(OH) nanoparticles and CNT f . An X-ray phase analysis of the samples was performed on a Rigaku SmartLab 3 diffractometer (Rigaku Corporation, Tokyo, Japan) with CuKα radiation. The diffraction (XRD) patterns were taken in the range of angles 2θ = 10-60 • at a speed of 1 • /min. The phase composition of the nanoparticles was determined with the use of PDWin 4.0 software (NPO "Burevestnik", St. Petersburg, Russia) using the profile analysis of XRD patterns. The results of the analysis were compared with the ASTM database. The crystallite sizes were calculated from the broadening of the X-ray diffraction lines according to the Scherrer equation [68].
A thermogravimetric analysis (TGA) of the materials was conducted using a DTG-60 setup (Shimadzu, Kyoto, Japan). The samples were heated up to 600 • C at a rate of 5 • C/min in air flow (100 mL/min). The thermal stability index, τ 10 (the temperature at which a polymer or a composite loses 10% of its initial weight because of thermal destruction) was determined using TGA curves.
A thermomechanical analysis (TMA) in extension mode was applied in order to investigate the behavior of the films under heating. A TMA 402 F1 Hyperion thermal Analyzer (NETZSCH, Selb, Germany) was used. The samples were heated at a rate of 5 • C/min in argon flow (70 mL/min). Glass transition temperatures T g were determined from TMA curves. A standard 0.5 MPa loading was applied to the PMDA-ODA-based samples. In the case of the films based on the flexible R-BAPS, the external stress was 25 kPa, since due to the steep transition to a plastic state, this polymer stretched far beyond the deformation range registered by the setup.
Mechanical characteristics of the nanocomposite and reference PI films were studied with the use of an AGS-X 5kN (Shimadzu, Kyoto, Japan) setup. Mechanical tests were performed in a uniaxial extension mode at room temperature. Strip-like samples 2 × 20 mm in size were stretched at a rate of 10 mm/min. The Young's modulus E, yield stress σ y , break stress σ b , and the ultimate deformation ε b for each sample were determined. For each composition 8-10 strips were tested, and the average values were calculated.

Structure and Morphology
The optical images of the surfaces of the nanocomposites we prepared are displayed in Figure 2. It can be seen that the clustering of the MO nanoparticles took place, because the micron-sized aggregates were registered. This indicates the certain heterogenization of the material. However, the results of the study of samples by SEM, as well as strong effects of the nanofiller on the functional characteristics of the nanocomposites we observed, of the material. However, the results of the study of samples by SEM, as well as strong effects of the nanofiller on the functional characteristics of the nanocomposites we observed, convincingly indicate that a large part of the nanofiller is uniformly distributed within the PI matrix. To investigate the interfacial interactions between the PI matrix and nanofillers and to assess the distribution of the nanoparticles within the matrix, we analyzed the SEM images ( Figure 3) of the fractured surfaces and AFM images ( Figure 4) of the surfaces of the films. The signs of plastic deformation were found on the fractured surfaces of all the nanocomposite samples (Figure 3b-f). It is apparent from Figure 3b that CNF had almost no adhesion to the matrix, providing a host of cavities inside the latter. In contrast, a cryo-cleavage of a CNTf-containing nanocomposite seemed uniform enough (Figure 3c). In the PI/TiO2 nanocomposite film, the nanoparticles were observed to aggregate, but they have good adhesion to the polymer matrix and are well-embedded in it (Figure 3d). Considering Figure 3e,f, one should conclude that the morphology of the nanocomposites filled with binary fillers is determined mostly by the carbon nanospecies, rather than by MO nanoparticles. Nonetheless, the sample with TiO2/CNF nanofiller has fewer voids compared with the single CNF-filled nanocomposite (Figure 3b,e). This forms the adhesion of the MO nanospecies to the matrix.    Figure 3b that CNF had almost no adhesion to the matrix, providing a host of cavities inside the latter. In contrast, a cryo-cleavage of a CNT f -containing nanocomposite seemed uniform enough (Figure 3c). In the PI/TiO 2 nanocomposite film, the nanoparticles were observed to aggregate, but they have good adhesion to the polymer matrix and are well-embedded in it (Figure 3d). Considering Figure 3e,f, one should conclude that the morphology of the nanocomposites filled with binary fillers is determined mostly by the carbon nanospecies, rather than by MO nanoparticles. Nonetheless, the sample with TiO 2 /CNF nanofiller has fewer voids compared with the single CNF-filled nanocomposite (Figure 3b,e). This results from the adhesion of the MO nanospecies to the matrix. of the material. However, the results of the study of samples by SEM, as well as strong effects of the nanofiller on the functional characteristics of the nanocomposites we observed, convincingly indicate that a large part of the nanofiller is uniformly distributed within the PI matrix. To investigate the interfacial interactions between the PI matrix and nanofillers and to assess the distribution of the nanoparticles within the matrix, we analyzed the SEM images ( Figure 3) of the fractured surfaces and AFM images (Figure 4) of the surfaces of the films. The signs of plastic deformation were found on the fractured surfaces of all the nanocomposite samples (Figure 3b-f). It is apparent from Figure 3b that CNF had almost no adhesion to the matrix, providing a host of cavities inside the latter. In contrast, a cryo-cleavage of a CNTf-containing nanocomposite seemed uniform enough (Figure 3c). In the PI/TiO2 nanocomposite film, the nanoparticles were observed to aggregate, but they have good adhesion to the polymer matrix and are well-embedded in it (Figure 3d). Considering Figure 3e,f, one should conclude that the morphology of the nanocomposites filled with binary fillers is determined mostly by the carbon nanospecies, rather than by MO nanoparticles. Nonetheless, the sample with TiO2/CNF nanofiller has fewer voids compared with the single CNF-filled nanocomposite (Figure 3b,e). This forms the adhesion of the MO nanospecies to the matrix.   nm which was likely caused by denser packing of the material containing TiO2 nan ticles. This corresponds well with SEM images (Figure 3b,e). Moreover, the MO particles on the film surface can evidently decrease the difference between the heig the topography maps. The values of surface roughness of CNTf-containing nanoco sites were quite close, Rq = 13.8 nm and 12.8 nm for R-BAPS/CNTf and R-BAPS/TiO2 compositions, respectively (Figure 4e,f). The FTIR spectrum of the pristine R-BAPS is shown in Figure 5a. The charact peaks of the imide rings at 1778 cm −1 , 1716 cm −1 (symmetric and antisymmetric stretching vibrations), 1373 cm −1 (C-N stretching vibrations) and 740 cm −1 (deform C=O vibrations) were registered [67]. The peaks at 1350 cm −1 (a shoulder near 1373 and 1147 cm −1 were ascribed to SO2-group vibrations. The spectra of R-BAPSnanocomposites (Figure 5b-e) also exhibited distinctive imide ring absorption p This indicates the completion of the PI imidization process not affected by the prese the fillers. However, some changes in the peak positions were observed. When CN CNF were inserted in the R-BAPS matrix, a certain shift to the low-frequency regio widening of the bands corresponding to symmetric and antisymmetric C=O stret vibrations of the imide cycle took place (Figure 5b), which may be attributed to th velopment of a network of hydrogen bonding. The shift of the same peaks R-BAPS/MO nanocomposites was rather small, which is likely due to the low conc tion of the nanoparticles. Filling R-BAPS with binary MO/nanocarbon mixtures led same changes in the shapes of the bands and peak positions of C=O stretching vibr as in the case of filling this PI with CNTf and CNF nanospecies individually (Figu   (Figure 4a) shows a smooth morphology with roughness R q = 3.8 nm (root mean square (RMS) roughness). We previously observed that morphology of PI-based nanocomposites is dependent on the type of MO [58]. The addition of TiO 2 to R-BAPS led to the increase in roughness up to R q = 6.3 nm (Figure 4b). At the same time, in the case of CNF (Figure 4c) as a filler, the surface roughness significantly increased up to R q = 32.5 nm. This can be explained by rather large sizes of CNF. The nanocomposite filled with a binary mixture of TiO 2 /CNF showed smoother morphology with R q = 24.6 nm which was likely caused by denser packing of the material containing TiO 2 nanoparticles. This corresponds well with SEM images (Figure 3b,e). Moreover, the MO nanoparticles on the film surface can evidently decrease the difference between the heights on the topography maps. The values of surface roughness of CNT f -containing nanocomposites were quite close, R q = 13.8 nm and 12.8 nm for R-BAPS/CNT f and R-BAPS/TiO 2 /CNT f compositions, respectively (Figure 4e,f).
The FTIR spectrum of the pristine R-BAPS is shown in Figure 5a. The characteristic peaks of the imide rings at 1778 cm −1 , 1716 cm −1 (symmetric and antisymmetric C=O stretching vibrations), 1373 cm −1 (C-N stretching vibrations) and 740 cm −1 (deformation C=O vibrations) were registered [67]. The peaks at 1350 cm −1 (a shoulder near 1373 cm −1 ) and 1147 cm −1 were ascribed to SO 2 -group vibrations. The spectra of R-BAPS-based nanocomposites (Figure 5b-e) also exhibited distinctive imide ring absorption peaks. This indicates the completion of the PI imidization process not affected by the presence of the fillers. However, some changes in the peak positions were observed. When CNT f and CNF were inserted in the R-BAPS matrix, a certain shift to the low-frequency region and widening of the bands corresponding to symmetric and antisymmetric C=O stretching vibrations of the imide cycle took place (Figure 5b), which may be attributed to the development of a network of hydrogen bonding. The shift of the same peaks in the R-BAPS/MO nanocomposites was rather small, which is likely due to the low concentration of the nanoparticles. Filling R-BAPS with binary MO/nanocarbon mixtures led to the same changes in the shapes of the bands and peak positions of C=O stretching vibrations as in the case of filling this PI with CNT f and CNF nanospecies individually (Figure 5c). The C-N and SO 2 -group stretching vibrations were also affected by the presence of the nanofillers. It is evident that ZrO 2(OH) nanoparticles augmented the intensity of the 1147 cm −1 band, the effect of TiO 2 on this band being very small (Figure 5d). Some changes in the intensities of the bands 1350 cm −1 and 1147 cm −1 along with a slight shift of the 1373 cm −1 peak to the low-frequency region were observed upon filling the R-BAPS with binary mixtures of MOs with nanocarbon ( Figure 5e). Overall, all the nanofillers are proved to affect the matrix due to the formation intermolecular bonds between the PIs' macromolecules and the surfaces of the nanoparticles. The nanosized TiO 2 , CNT f , and CNF interact both with the imide cycle (this presumably results in the changes in packing of the macromolecules) and with SO 2 -groups, while ZrO 2(OH) nanoparticles connect predominantly with SO 2 -groups. The C-N and SO2-group stretching vibrations were also affected by the presence of the nanofillers. It is evident that ZrO2(OH) nanoparticles augmented the intensity of the 1147 cm −1 band, the effect of TiO2 on this band being very small (Figure 5d). Some changes in the intensities of the bands 1350 cm −1 and 1147 cm −1 along with a slight shift of the 1373 cm −1 peak to the low-frequency region were observed upon filling the R-BAPS with binary mixtures of MOs with nanocarbon ( Figure 5e). Overall, all the nanofillers are proved to affect the matrix due to the formation intermolecular bonds between the PIs' macromolecules and the surfaces of the nanoparticles. The nanosized TiO2, CNTf, and CNF interact both with the imide cycle (this presumably results in the changes in € packing of the macromolecules) and with SO2-groups, while ZrO2(OH) nanoparticles connect predominantly with SO2-groups.      The positive influence of titania and zirconia nanoparticles is usually attributed to a certain degree of linking between PI macrochains and the surfaces of the nanoparticles via hydrogen bonding. One should mention that this effect depends strongly on the size, concentration, and surface functionality of MO nanospecies [50]. For instance, the smallest nanoparticles have the highest surface energy and a tendency to aggregate. In this case, the aggregates become stress concentrators and loosen the polymer structure (compare R-BAPS/ZrO2(18nm) with R-BAPS/ZrO2(8nm) samples in Figure 8). Mechanical characteristics of PI/MO nanocomposites are also affected by the surface functionality of the nanoparticles. The stiffness values for PI/ZrO2(OH) nanocomposites were observed to be lower than those for the pristine PIs. This implies a negative effect of surface OH-groups on the structure of the PI material. Figure S3a,c shows that ZrO2(OH) nanoparticles form large aggregates dispersed ununiformly within the matrix, a lot of cavities  influence of titania and zirconia nanoparticles is usually attributed to a certain degree of linking between PI macrochains and the surfaces of the nanoparticles via hydrogen bonding. One should mention that this effect depends strongly on the size, concentration, and surface functionality of MO nanospecies [50]. For instance, the smallest nanoparticles have the highest surface energy and a tendency to aggregate. In this case, the aggregates become stress concentrators and loosen the polymer structure (compare R-BAPS/ZrO 2 (18nm) with R-BAPS/ZrO 2 (8nm) samples in Figure 8). Mechanical characteristics of PI/MO nanocomposites are also affected by the surface functionality of the nanoparticles. The stiffness values for PI/ZrO 2(OH) nanocomposites were observed to be lower than those for the pristine PIs. This implies a negative effect of surface OH-groups on the structure of the PI material. Figure S3a,c shows that ZrO 2(OH) nanoparticles form large aggregates dispersed ununiformly within the matrix, a lot of cavities being also formed. This changes the morphology and topography of the PI, substantially increasing its roughness. The more palpable reinforcing effect of the nanocarbon species is ensured by the formation of a strong network consisting of 1D particles. This network is capable of taking over the external mechanical load from a polymer matrix [69,70]. Moreover, with CNT f being preliminarily functionalized, chemical interactions between the PI macromolecules and these nanoparticles are provided, imparting even better properties to such nanocomposites [61].  The addition of the binary nanofillers to the PIs led to an increase in their rigidity, the Young's moduli turning out to become higher than those of the PIs filled with only MO nanoparticles. For instance, the elasticity modulus of the PMDA-ODA/TiO 2 /CNF and PMDA-ODA/TiO 2 /CNT f increased by 11% regarding the corresponding PI/TiO 2 composite. It is worth noting that the difference between these values was more sensible if PI/MO nanocomposites had originally poor properties, e.g., PI/ZrO 2(OH) and PI/ZrO 2 (8nm) samples vs. PI/ZrO 2(OH) /nanocarbon and PI/ZrO 2 (8nm)/nanocarbon composites. For example, the Young's modulus of the R-BAPS/ZrO 2(OH) /CNF sample was 54% higher than that of the R-BAPS/ZrO 2(OH) nanocomposite (Figure 8a). It is the nanocarbon networks that determine the stiffness and mechanical strength of the nanocomposites, with the MO nanoparticles being embedded and bridged into the network by 1D carbon nanospecies. The differences in the morphology and topography of the nanocomposites containing single ZrO 2(OH) filler and binary ZrO 2(OH) /CNF can be seen in Figure S3. There are two groups of mechanical characteristics reflecting the behavior of the film polymer material at different stages of the deformation process. Each of the groups is affected by different factors. For instance, the values of E and σy are determined predominantly by the strength of the system of intermolecular bonds in the material. On the other hand, break stress and elongation at break reflect the features of the material's morphology, such as structural heterogeneity, and the presence and concentration of local internal defects, which occur at the phase boundaries during the processing of the nanocomposite. The differences in the magnitude of the characteristics from the second group (σb and εb) show exactly the differences in the degree of microheterogeneity of the structure of the nanocomposite materials formed by adding various nanoparticles into the polyimide matrix. However, one should mention that from a practical viewpoint, such characteristics as E and σy are of great importance since generally PI materials are exploited exactly in the range of small deformations where the behavior of a material is characterized by these two parameters.

Thermal Properties
The thermal stability indices τ10 of the pristine PIs and PI-based nanocomposites are shown in Figure 9. All the materials are proven to possess excellent thermal stability. The thermal degradation of nanoparticles per se (both metal oxides and carbon nanoparticles) begins at temperatures significantly higher than the onset temperatures of thermal degradation of the matrix polymers studied in the work. Thus, the influence of nanosized fillers on the thermal stability indices of the nanocomposite materials determined in the There are two groups of mechanical characteristics reflecting the behavior of the film polymer material at different stages of the deformation process. Each of the groups is affected by different factors. For instance, the values of E and σ y are determined predominantly by the strength of the system of intermolecular bonds in the material. On the other hand, break stress and elongation at break reflect the features of the material's morphology, such as structural heterogeneity, and the presence and concentration of local internal defects, which occur at the phase boundaries during the processing of the nanocomposite. The differences in the magnitude of the characteristics from the second group (σ b and ε b ) show exactly the differences in the degree of microheterogeneity of the structure of the nanocomposite materials formed by adding various nanoparticles into the polyimide matrix. However, one should mention that from a practical viewpoint, such characteristics as E and σ y are of great importance since generally PI materials are exploited exactly in the range of small deformations where the behavior of a material is characterized by these two parameters.

Thermal Properties
The thermal stability indices τ 10 of the pristine PIs and PI-based nanocomposites are shown in Figure 9. All the materials are proven to possess excellent thermal stability. The thermal degradation of nanoparticles per se (both metal oxides and carbon nanoparticles) begins at temperatures significantly higher than the onset temperatures of thermal degradation of the matrix polymers studied in the work. Thus, the influence of nanosized fillers on the thermal stability indices of the nanocomposite materials determined in the work is ensured by their impact on the system of intermolecular bonds and the morphology of the material (in particular, on the structure of the interfaces between the polymer and fillers). It is obvious from Figure 9 that nanoparticles affect the thermal stability of the PIs variously. As we already discussed in [54,58] the effect is determined by a number of factors including the type, size, surface functionality, and concentration of nanofiller. Considering MO nanoparticles, their positive influence on the thermal stability of PIs containing SO 2 -groups (R-BAPS) can be caused by two reasons: The first one is presumed to be the ability of these nanospecies to take part in chemical reactions with atmospheric oxygen at elevated temperatures providing additional links between positively charged S atom in PI radicals and the surface of nanoparticles bearing negative charge due to active oxygen species (AOS) formed. These links can slow down the degradation of a nanocomposite film in a certain region of temperatures (Figure 9a). However, the AOS may well have a detrimental effect on a PI matrix (PMDA-ODA) with no sulfonyl groups catalyzing its thermo-oxidative decomposition (Figure 9b), e.g., in PMDA-ODA/3%ZrO 2(OH) nanocomposite [23,58]. On the other hand, improved mechanical properties of PI/MO nanocomposites discussed above imply good interfacial interactions between the matrix and filler. These interactions can provide a physical barrier, retarding the out-diffusion of volatile products during thermal degradation of the nanocomposites. The effect of CNF on the τ 10 values of PIs is generally less pronounced, at least at the concentrations provided. This agrees well with the data we obtained earlier [54]. The difference in thermal behavior of CNF-and CNT f -containing nanocomposites can be attributed to the profound contrast in their morphology (Figure 3b,c). The denser and more homogeneous structure of R-BAPS/CNT f nanocomposite is evidently responsible for its high thermal stability.  Figures 10a and 11a illustrate TMA curves of a series of the PMDA-ODA-and R-BAPS-based samples. The nanofillers were demonstrated to affect Tg diversely (Figures  10b and 11b). Almost all the MO nanoparticles generally decreased Tg of both the PIs. We discussed such an effect in detail in [58]. It is noteworthy that the influence of the MO nanospecies on Tg values is more substantial in PMDA-ODA samples. This can be explained by the fact that the flexible macromolecules of R-BAPS facilitate its segmental mobility at a temperature Tg lower than that of PMDA-ODA. Apparently, the structure of the polymer rather than the nanofiller plays a decisive role in the behavior of R-BAPS-based nanocomposites in the vicinity of the Tg. It is also evident from Figure 10b that two types of nanoparticles augment the Tg of PMDA-ODA, namely ZrO2(OH) and CNTf, whose surfaces bear functional groups. One can presume additional interactions between these groups and PMDA-ODA macromolecules retarding the segmental mobility of the polymer and increasing its Tg.

Thermomechanical Properties
It is seen from Figure 10a that the pristine PMDA-ODA was stretched quite strongly The introduction of binary nanofillers into the PI matrices causes the enhancement of their thermal stability, which becomes predominantly even higher than those of the corresponding single-component nanocomposites. The most pronounced effect was observed in the nanocomposites filled with the TiO 2 /CNT f combination, whose τ 10 were augmented by more than 30 • C compared with these values for the pristine R-BAPS and PMDA-ODA matrices.
Putting together SEM images ( Figure 3) and the data on mechanical characteristics (Figure 8), the synergism in R-BAPS-based nanocomposites with binary nanofillers is supposed to stem from the simultaneous change in the structure of the matrix (provided by carbon nanospecies) and specific interactions of the MO nanoparticles with SO 2 -containing macrochains. For instance, comparing Figure 3d,e one can notice that the morphology of a nanocomposite with mixed TiO 2 /CNT f differs considerably from that of the sample with the single TiO 2 nanofiller and is determined by the nanocarbon filler. It is seen from Figure 3b, that the R-BAPS/CNF sample has many caverns due to the presence of CNF. These caverns may well-facilitate the out-diffusion of the PI degradation products. Despite the R-BAPS/TiO 2 /CNF structure being loose, one can observe an increase in thermal stability of the nanocomposite, which is supposed to be ensured by titania. As for PMDA-ODA-based samples, the positive effect of the binary filler on their thermal stability obviously results primarily from the change in the structure of the materials due to the addition of the nanocarbon. Moreover, carbon nanospecies can likely cover the active surface of the MO nanoparticles (Figure 2b,c), thereby hindering their catalytic effect on thermo-oxidative destruction of PMDA-ODA.
Similar results were obtained in [54] for PMDA-ODA and another PI-bearing sulfonyl groups (DPhO-BAPS) filled with binary CeO 2 /nanocarbon mixtures. The synergistic effect of a mixture of carbon nanoparticles (GNPs) with inorganic salt (BN) on thermal stability of a PI was also registered in [52], but the concentration of the salt reached 30 wt.% so as to provide an increase in τ 5 value less than 10 • C.

Thermomechanical Properties
Figures 10a and 11a illustrate TMA curves of a series of the PMDA-ODA-and R-BAPSbased samples. The nanofillers were demonstrated to affect T g diversely (Figures 10b and 11b). Almost all the MO nanoparticles generally decreased T g of both the PIs. We discussed such an effect in detail in [58]. It is noteworthy that the influence of the MO nanospecies on T g values is more substantial in PMDA-ODA samples. This can be explained by the fact that the flexible macromolecules of R-BAPS facilitate its segmental mobility at a temperature T g lower than that of PMDA-ODA. Apparently, the structure of the polymer rather than the nanofiller plays a decisive role in the behavior of R-BAPS-based nanocomposites in the vicinity of the T g . It is also evident from Figure 10b that two types of nanoparticles augment the T g of PMDA-ODA, namely ZrO 2(OH) and CNT f , whose surfaces bear functional groups. One can presume additional interactions between these groups and PMDA-ODA macromolecules retarding the segmental mobility of the polymer and increasing its T g .
It is seen from Figure 10a that the pristine PMDA-ODA was stretched quite strongly above the T g . When the temperature reached~440 • C the deformation rate decreased, signaling an initial stage of the polymer thermal degradation accompanied with the formation of destruction crosslinks. The crosslinks are likely to stiffen the material in this temperature region [67,71]. To assess the rigidity of the samples above the T g , we determined a value of deformation in the temperature range between the T g and a temperature of maximum deformation on the TMA curves, ∆ε (Figure 10c). Obviously, MO nanoparticles affect the ∆ε value of the nanocomposites diversely, the rigidity of the materials depending on the type, size, and concentration of the nanospecies. Generally, the increase in ∆ε value is observed when either the size or concentration of the MO nanoparticles is augmented, since more aggregates preventing polymer-polymer interactions are formed and more flexible behavior in PMDA-ODA-based nanocomposites is demonstrated after glass transition [58]. As concerns the nanocomposites filled with nanocarbon (either CNT f or CNF), the ∆ε value was registered to drop against the host PMDA-ODA. The network of 1D carbon nanoparticles is presumed to contribute to the crosslinking at elevated temperatures and impede the motion of macromolecular segments [54]. Filling PMDA-ODA with binary MO/nanocarbon mixtures resulted in a rise in its stiffness compared with the corresponding nanocomposites with single MO nanofillers, implying that it is the nanocarbon network that ensures low ∆ε values of such materials. Unlike PMDA-ODA, flexible R-BAPS stretches quite well even under little stress applied, exhibiting elastic behavior above the T g followed by plastic deformations at higher temperatures. The mutual motion of the macrochains as a whole is eased significantly at temperatures ca. 50 • C higher than the T g . As a result, this PI begins flowing (temperature at which the flowing occurs is denoted as T fl ) (Figure 11a). Such behavior is typical of the PIs containing several bridge groups [54,72]. Analyzing Figure 11c, one can observe that the MO nanospecies strongly affect T fl decreasing this value by more than 25 • C. The quasi-spherical MO nanoparticles plausibly hinder the polymer-polymer interactions, facilitating a relative motion of the R-BAPS macromolecules. The same effect can be responsible for a drop in the T fl in R-BAPS/CNT f sample, since the polymer-nanoparticle (either MO or CNT f ) crosslinked structures determining the stiffening of the corresponding nanocomposites at room temperature ( Figure 8) may well be broken or at least become more flexible at higher temperatures. Despite the fact that CNF makes the polymer packing less dense (Figure 3b,e), which should promote the motion of the macrochains, the T fl of the R-BAPS nanocomposites containing CNF is much higher than that of the pristine matrix. It appears that the mobility of the sample as a whole can be hindered because of a rigid network formed by large CNF incapable of "flowing" until a high temperature is reached. This network obviously determines the behavior of the R-BAPS-based nanocomposites filled with binary MO/CNF nanofillers, since their T fl always surpasses the T fl of the unfilled polymer.

Conclusions
In this research, we reported on the fabrication technique and comprehensive study of the thermal and mechanical properties of PI-based nanocomposites filled with binary mixtures of MO nanoparticles (either TiO2 or ZrO2) and nanocarbon (either CNF or CNTf). The structural and morphological features of the aforesaid materials were considered as well. The improvement of mechanical characteristics compared with single-filler (MO nanospecies) nanocomposites was demonstrated. It is the nanocarbon additive that is supposed to be responsible for the increase in stiffness (both below and above Tg) and yield stress of the materials. On the other hand, the MO nanoparticles are likely to augment the thermal stability of the PI. As a result, novel materials with an extended set of enhanced properties were developed. The most pronounced synergistic effect of the constituents was registered in PIs filled with the binary TiO2/CNTf mixture whose thermal stability (τ10 value) turned out to be higher than both single-filler nanocomposites, the films possessing excellent mechanical properties.
Filling PI with binary MO/carbon nanoadditives was proven to be an effective method of fabrication of nanocomposites to overcome the limitations of the single-filler systems. We believe that such a versatile approach to the design of new nanocomposite materials would be an ideal choice in high-end engineering applications.

Conclusions
In this research, we reported on the fabrication technique and comprehensive study of the thermal and mechanical properties of PI-based nanocomposites filled with binary mixtures of MO nanoparticles (either TiO 2 or ZrO 2 ) and nanocarbon (either CNF or CNT f ). The structural and morphological features of the aforesaid materials were considered as well. The improvement of mechanical characteristics compared with single-filler (MO nanospecies) nanocomposites was demonstrated. It is the nanocarbon additive that is supposed to be responsible for the increase in stiffness (both below and above T g ) and yield stress of the materials. On the other hand, the MO nanoparticles are likely to augment the thermal stability of the PI. As a result, novel materials with an extended set of enhanced properties were developed. The most pronounced synergistic effect of the constituents was registered in PIs filled with the binary TiO 2 /CNT f mixture whose thermal stability (τ 10 value) turned out to be higher than both single-filler nanocomposites, the films possessing excellent mechanical properties.
Filling PI with binary MO/carbon nanoadditives was proven to be an effective method of fabrication of nanocomposites to overcome the limitations of the single-filler systems.
We believe that such a versatile approach to the design of new nanocomposite materials would be an ideal choice in high-end engineering applications.